absorption and diffusion of hydrogen in steels

H. J. GRABKE, E. RIECKE: ABSORPTION AND DIFFUSION OF HYDROGEN IN STEELS

ABSORPTION AND DIFFUSION OF HYDROGEN INSTEELS

ABSORPCIJA IN DIFUZIJA VODIKA V JEKLIH

Hans Jürgen Grabke, Ernst RieckeMax-Planck-Institut für Eisenforschung GmbH, Postfach 140444, 40074 Düsseldorf, Germany

Prejem rokopisa - received: 2000-12-11; sprejem za objavo - accepted for publication: 2000-12-20

To understand and control hydrogen-induced cracking and hydrogen-induced stress-corrosion cracking in steels, thefundamental processes of hydrogen absorption and diffusion are of interest.The electrochemical hydrogen permeation method has been extensively applied to obtain data on hydrogen absorption anddiffusion in iron, iron alloys and steels. Steady-state permeation measurements yield the permeation coefficient and thehydrogen concentration and activity aH, which are established on the surface in electrolytes. This information was used to studythe effects of alloying elements on aH present during corrosion.Permeation transients from non-steady-state measurements yield the hydrogen diffusivity, which is affected by the presence ofhydrogen traps. For shallow traps with low binding energies a distribution equilibrium between trap and lattice sites applies,deep traps are already saturated at low aH and hydrogen in deep traps is immobile at ambient temperature. The effects of Mo, V,Nb, Ti, Zr and their carbides and nitrides on the hydrogen diffusivity were studied and information was derived on the numbersand binding energies of traps which are caused by strain fields around big atoms of alloying elements or in the neighbourhoodof their precipitates.For some pipeline steels with different microstructures, and after welding simulation, hydrogen diffusivities and trapping werestudied. The results were related to hydrogen-induced stress-corrosion cracking, which was observed in constant-extension-ratetests. The work of fracture decreased with increasing hydrogen activity, i.e. the mobile hydrogen is decisive for failure, and thehydrogen in deep traps has no obvious effect on the fracture behaviour of steels.

Key words: iron alloys, pipeline steels, hydrogen activity, permeation, diffusivity, fracture energy

Za razumevanje in kontrolo pokanja zaradi vodika in napetostnega korozijskega pokanja jekel je treba poznati temeljne proceseabsorpcije in difuzije vodika.Metoda elektrokemi~ne permeacije vodika se je mnogo uporabljala za pridobivanje podatkov o absorpciji in difuziji vodika vèlezu, njegovih zlitinah in jeklih. Iz meritev permeacije v stacionarnem stanju so bili dolo~eni koeficient permeacije terkoncentracija in aktivnost vodika na povr{ini kovine v elektrolitih. Metoda je bila tudi uporabljena za dolo~anje vpliva legirnihelementov na aktivnost vodika med korozijo.Tranzienti permeacije pri nestacionarnih meritvah dajo podatke o difuzivnosti vodika, na katero vplivajo prisotne pasti. Zaplitve pasti z majhno vezno energijo se vzpostavi ravnote`je med pastmi in mre`nimi mesti.Globoke pasti so nasi~ene è pri majhni aktivnosti vodika, vodik v njih pa je nepremi~en pri temperaturi okolice. Raziskani sobili vplivi Mo, V, Nb, Ti, Zr ter njihovih karbidov in nitridov na difuzivnost vodika in pridobljeni podatki o {tevilu in veznienergiji pasti. Te nastanejo zaradi napetostnih polj okoli velikih atomov legirnih elementov in v okolici njihovih precipitatov.Za nekatera jekla za cevovode z razli~no mikrostrukturo, tudi po simulaciji varjenja, sta bila raziskana difuzivnost in ujetjevodika v pasti, povezana z vodikom, induciranim napetostnim korozijskim pokanjem, ki je bilo opredeljeno z nateznimipreizkusi s konstantno hitrostjo iztezanja. Delo preloma se zmanj{uje z ve~anjem aktivnosti vodika, kar pove, da je mobilnivodik odgovoren za prelom, vodik v globokih pasteh pa ne vpliva na prelomno vedenje jekel.

Klju~ne besede: èlezove zlitine, jekla za cevovode, vodik, aktivnost, permeabilnost, difuzivnost, energija preloma

1 INTRODUCTION

Hydrogen in iron is an important and intensivelystudied system 1-11. Severe failure can be caused by thehydrogen embrittlement of steels, which appears as:surface blistering; hydrogen-induced cracking (HIC);stress-corrosion cracking (HISCC); classical, weld andstress cracking and hydrogen-environment embrittlement(HEE).

After adsorption and dissociation hydrogen isdissolved from the atmosphere by absorption accordingto:

H2 (gas) = 2H (dissolved) (1)

and its concentration follows ’Sieverts law’

cH = KS ⋅ (pH2)1/2 (2)

where cH [mol/cm3] is the concentration of dissolved Hatoms, KS is Sieverts constant and pH2 [bar] thehydrogen pressure in the environment. The solubility at1 bar H2

10-14 is extremely low at room temperature, just3⋅10-6 at %, but it rises with temperature to 1.6⋅10-2 at%at 900°C in α-Fe. This corresponds to 3cm3 of H2 per100g Fe. In γ-Fe the solubility is a little higher, 2.3⋅10-2

at% at 900 °C and in the melt cH is much higher, 0.13at% at 1540 °C. As a consequence it is dangerous tosaturate steels with hydrogen at elevated temperatures,without the possibility for H2-desorption during cooling.This desorption may be hindered by oxide scales,enamel or zinc coatings; and an oversaturated solutionmay result in a variety of failures: shatter cracks, flakes,fish eyes, etc.

MATERIALI IN TEHNOLOGIJE 34 (2000) 6 331

UDK 669.14.018.85:669.778:620.193 ISSN 1580-2949Izvirni znanstveni ~lanek MATER. TEHNOL. 34(6)331(2000)

Hydrogen absorption occurs not only from gaseoushydrogen, but more often from electrolytes, e.g. duringpickling or the corrosion of steels. The most importantreactions during the corrosion of iron in an acidenvironment are 15-18:the anodic dissolution of Fe (Volmer reaction):

Fe + 2 H+ = Fe2+ + 2 Had (3a)

the recombination of adsorbed H (Tafel reaction):

2 Had = H2 (gas) (3b)

and the absorption of Had:

Had = H(dissolved) (3c)

The interplay of these reactions establishes ahydrogen activity aH on the metal surface which isstrongly affected by inhibitors and promoters in theelectrolyte and by the presence of alloying elements andimpurities in the metal surface 19-25. The study of theselatter effects was the subject of our studies 26-28. Thehydrogen activity is defined as the hydrogenconcentration cH related to the hydrogen concentrationcH

0 at T = 0°C and at 1 bar H2

ac

c

p

pH

H

HO

H

HO

2

2

= =

1 2/

(4)

If the anodic dissolution reaction (3a) is enhanced(e.g. by the presence of sulfur) and/or the recombinationreaction (3b) is retarded (e.g. by arsenate) very highhydrogen activities can occur on steel surfaces whichcause the severe absorption of hydrogen.

Hydrogen is very dangerous, especially in ferriticsteels, since it generally diffuses very quickly to defects,where it exerts its deleterious effects. In fact, thediffusivity of H in pure and undeformed α-iron at room

temperature is very high, about 1⋅10-4 cm2/sec 27-32, thiscompares with about 1⋅10-16 cm2/sec for carbon andnitrogen.

Our own careful studies on very pure iron yielded 6,7:

DH = (5.12 ± 0.6) ⋅ 10-4 exp (-(4.15 ± 0.3)kJ mol-1/RT)(5)

The purer and better recrystallized the iron samplesthe higher were the data obtained, see Figure 1. Theactivation energy is strikingly low: in comparison withabout 80 kJ/mol for carbon and nitrogen. This indicatesthat hydrogen moves through the ferritic lattice as aproton, tunneling through the interstitial sites. Neithergrain boundary nor surface diffusion are faster processes.Alloying elements and deformation generally decreasethe diffusivity of H in iron and a tremendous number ofArrhenius lines have been published over the years, onhydrogen diffusion in various steels, see Figure 2 6,7,33.

The data in Figure 2 were derived using theassumption that hydrogen concentration and diffusionfollow Fick’s laws, and for the higher temperature rangethe data obtained could be approximately extrapolated toequ. (5), whereas the data at lower temperatures aremostly on lower Arrhenius lines and would yield higheractivation energies. This diffusion behaviour ofhydrogen was treated by McNabb and Foster 33 byapplying a trapping model: ie. the hydrogen atoms


332 MATERIALI IN TEHNOLOGIJE 34 (2000) 6

Figure 1: Arrhenius plot of hydrogen diffusion coefficient (DH) inpure annealed iron, own data 6, 7 and nearby Arrhenius lines fromother authors 29-32

Slika 1: Arrheniusov diagram odvisnosti koeficienta difuzije DH v~istem àrjenem èlezu, lastni podatki 6,7 in pribli`ne Arrheniusove~rte za druge avtorje 29-52

Figure 2: Arrhenius plot of hydrogen diffusivities (Deff) measured bynumerous authors in different steels, compared to DH in purerecrystallized iron, and the H diffusivity in 60% cold-rolled ironwhich can be described by equation (12) with Nft = 6.7⋅1019/cm3Feand ∆H°ft = -36 kJ/molH 6, 7

Slika 2: Arrheniusov diagram odvisnosti koeficienta difuzije DH pomeritvah mnogih avtorjev v razli~nih jeklih v primerjavi z DH v~istem rekristaliziranem èlezu in difuzivnost H v 60 % hladnovaljanem jeklu, ki jo opisuje ena~ba (12) z Nft = 6,7⋅1019/cm3Fe in∆H°ft = -36 kJ/molH 6, 7

wander in a random manner through the lattice, but tendto get trapped or delayed at sites which act as ’traps’.These sites are regarded as potential wells of greaterdepth, than those in the undisturbed, defect-free crystallattice. During the hydrogen diffusion, a localequilibrium is assumed for the ’trapping reaction’:

H(in normal, interstitial site) + trap == H(in trap) + free normal, interstitial site (6)

This situation corresponds to an adsorptionequilibrium where:

rate of adsorption = rate of desorption:k⋅cH (1-θ)⋅Nt = k'⋅ct⋅N (7)

where Nt is the number of traps and θ is their degree ofoccupation, N is the number of normal interstitial sitesand ct = θ⋅Nt is the concentration of H in the traps. Thequotient of the rate constants for the forward andbackward reactions is the equilibrium constant of thetrapping reaction (6):

k

k'

N

c (1 )H

= ⋅−

=θθ

K = exp (-∆Gt0 / RT)

= exp (+∆St0 / R) exp (-∆Ht

0 / RT) (8)

and for the degree of trap occupation:

θ ββt =

⋅+

=+

c K / N

1 c K / NH

H 1β = cH ⋅ K/N (9)

In steels there are different kinds of traps, in principlea more-or-less continuous spectrum of ’binding energies’∆Ht° and values of K is to be expected, however, weusually distinguish between flat traps with-∆Hft < 30 kJ/molH and β << 1 and deep traps with-∆Hdt > 50 kJ/molH and β >> 1

The deep traps are already saturated at low values ofaH and cH and do not affect the diffusivity DH, whereasthe occupation of the flat traps varies with the increaseand the decrease of aH resp. cH in diffusion experimentsor during corrosion.

Non-steady-state diffusion experiments are generallydescribed by Fick’s second law, which in the presence oftraps describes the H concentration changes in thenormal interstitial sites and in the flat traps:

∂∂

∂∂

∂∂

c

t

c

tD

c

xH ft

HH

2+ =

2

(10)

this equation can be rewritten:

∂∂ ∂ ∂

∂∂

c

t

D

c c

c

xH H

ft H

H

2=

+

1

2

( / )(11)

A comparison with the usual formulation of Fick’ssecond law shows that in the presence of flat traps aneffective diffusivity of H is observed, which is given by

DD

c c

Deff

H

ft H

H=+

=+1 1( / )∂ ∂ α

(12)

Since cft = θ ⋅ Nft we obtain

α = Kft ⋅ Nft / N= (Nft / N) exp (∆Sft

0 / R) exp (-∆Hft0 / RT) (13)

Thus, hydrogen diffusivity decreases with a highoccupancy of flat traps Nft and high binding energiesI∆H°I, and the Arrhenius plots, as shown in Figure 2,can be explained by the trapping model and equation(12). Figure 2 also demonstrates that the numerousstudies on different steels might well be explained by theeffects of strained regions and dislocations 6,7. Thehydrogen diffusivity in iron which has been 60%deformed by cold-rolling corresponds to the diffusionbehaviour in many different steels with various degreesof deformation. It is probable that the flat trapscorrespond to dilated sites near dislocations and the deeptraps to dislocation cores. Such dilated sites are alsocaused by large atoms of alloying elements and by thestresses around precipitates.

Another aim of this study was to clarify the effects ofalloying elements such as Ti, Zr, V, Nb, Cr and Mo andtheir carbides and nitrides on hydrogen solubility anddiffusivity, to obtain a better understanding of traps.Studies were conducted on H permeation, diffusion andsolubility in several binary alloys of iron and the abovementioned elements, see table 1 34-40.

2 EXPERIMENTAL

2.1 Electrochemical Double Cell

Hydrogen absorption, permeation and diffusion weremeasured using the electrochemical double cell, firstdescribed by Devanathan and Stachurski 41,42. An iron orsteel membrane is charged with hydrogen from the entryside, at the exit side the hydrogen is oxidized anodicallyand the anodic current Ip is a measure of the hydrogenpermeation. The schematics in Figure 3a demonstratethe principle of the method. Two techniques forhydrogen charging have been applied a) electrolyticcharging from an aqueous solution or b) gas-phasecharging from pure H2 or N2-H2 mixtures. Theapplication of both techniques us allows to measure atlow and high hydrogen activities and at a very welldefined standard activity aH = 1, corresponding to pH2 =1 bar.

For the study of the compositional or structuraleffects of H absorption, the entry side remaineduncoated; this was also the case for sufficiently highcathodic polarisation currents, obtained at constantcathodic current densities iC or cathodic potentials EC.Otherwise, especially in the case of hydrogen absorptionfrom the gas phase, the entry side of the membrane hadto be Pd coated, to avoid the inhibition of hydrogenadsorption and absorption by oxide films. The exit sideof the membrane was always coated with Pd, andpolarised anodically at Ea = 200 mV(SHE) in a 0,1N



NaOH solution, deaerated by bubbling with N2. Here thehydrogen is oxidized immediately and its concentrationis virtually zero. The anodic permeation current (Ip) iscorrected by a blind current, measured when nohydrogen was charged.

2.2 Determination of the permeation coefficient

The permeation coefficient is the product of thehydrogen solubility at 1 bar H2 and the hydrogendiffusivity in a steady-state experiment:

Φ0 = ⋅c DH0

H (14)

Its value was determined for each materialinvestigated by charging the Pd-coated membrane fromthe atmosphere, i.e. pure hydrogen at 1 bar. Afterattaining a steady state, the flow of hydrogen (jH) isgiven by Fick’s first law:

j Dc c

dD c / d / dH H

H0

He

H H0=

−= ⋅ = Φ0 (15)

where d is the membrane thickness; the hydrogenconcentration at the entry side (cH°) and at the exit side(cHe) = 0. The anodic oxidation of the flux of hydrogenarriving at the exit side yields the permeation current:

Ip = jH ⋅ F ⋅ A (16)

where F is the Faraday constant and A is the area of themembrane. The permeability results from:

Φ0 =⋅⋅

= ⋅

I d

F A

mol

cmsecc D

p

H0

H (17)

2.3 Measurement of hydrogen activity

In electrolytes, hydrogen activitiy on iron and steelsurfaces is the result of an interplay of electrochemicalreactions (3). The value of aH can be very high due to thepromotion of reaction (3a), the discharge and Hadsorption, or to the inhibition of reaction (3b) - the Hrecombination (Tafel) reaction. It is of great scientificand technical interest to know about effects of theelectrolyte and of the steel composition on theestablished hydrogen activity, e.g. during pickling orcorrosion. In this study, permeation measurements wereused to obtain the hydrogen activity (aH) duringcorrosion in 1 M H2SO4.

The hydrogen concentration (cH) at the entry side ofthe corroding membrane is determined using:

cI d

F A c DH

p

H0

H

=⋅

⋅ ⋅ ⋅(18)

and related to the hydrogen concentration cH° at aH = 1(pH2 = 1 bar) by:

ac

cH

H

H0

= =I d

F A c D

p

H0

H

⋅

⋅ ⋅ ⋅=

I d

F A

p ⋅

⋅ ⋅Φ0(19)

Thus, the hydrogen activity can be obtained if thepermeation coefficient Φ° is known from the experimentdescribed above, and Φ° must be determined separatelyfor each material, since its value is dependent on thesteel composition and microstructure.

2.4 Determination of Hydrogen Diffusivity

Upon charging or discharging the membrane in theelectrochemical double cell, current transients areobserved (see Figure 3b). These transients correspond tothe absorption of H into the material during charging andthe emptying by anodic oxidation during discharging.The process does not only involve the hydrogen in idealsolution, i.e. in the normal interstitial lattice sites, butalso the flat traps are filled and emptied, whereas thedeep traps are saturated in the very first charging step(see Figure 3b) and H sticks to these traps during



Figure 3: Schematics of measurements in the electrochemical doublecell: a) schematic drawing of the double cell, hydrogen charging fromthe left side, from electrolyte or gas phase, anodic oxidation of thepermeating hydrogen in the right cell - measurement of the permeationcurrent IP; b) schematic presentation of a permeation experiment, inthe first step saturation of the deep traps at low charging current iC andlow hydrogen activity aH, then transients of hydrogen permeation fordetermination of the effective diffusivity Deff

Slika 3: Shema meritev v dvojni elektrokemi~ni celici: a) Shemadvojne celice, nasi~enje z vodikom z leve strani iz elektrolita ali izplina, anodna oksidacija permeatiranega vodika v desni celici -meritev toka permeacije Ip; b) Shema poskusa permeacije, v prvistopnji nasi~enje globokih pasti pri majhnem toku ic in majhniaktivnosti vodika aH, nato tranzienti permeacije vodika za dolo~itevdejanske difuzivnosti Deff

long-term discharging. In our experiments fordetermining the hydrogen diffusivity the deep traps werealways saturated in the first charging step at a lowcharging current (iC) and hydrogen activity (aH) and thenthe transients of charging and discharging weremeasured. These non-steady-state experiments lead to aDH for pure and undeformed iron, and according toequation (12), to a Deff, which is affected by the flattraps. Such a measurement is demonstrated in Figure 3b,where the permeation current during a sequence ofgalvanostatically controlled charging steps is plottedschematically. In the first step at (iC)O a low hydrogenactivity is established on and in the material and all thedeep traps are saturated. Then the cathodic current israised and lowered at a higher level, in the range of iCand aH, where the flat traps can be occupied and released.The permeation transients obtained in that range aredescribed by a series

I (t) I

I (0) In

p pS

p pS

n

n

−

−= − −−

=

∞

∑2 1 1

1

2 2( ) exp( )π τ (20)

where τ = Deff · t / d2. For the determination of Deff fromthe breakthrough time (tb) (see Figure 3b), the firstthree terms of the right hand side of equation (20) areused to give:

Deff ≈ d2 / 20 · tb (21)

The breakthrough time is defined by the intersectionpoint of the time axis with the tangent drawn at theinflection point of the permeation curve (see Figure 3b).

In the first charging of a hydrogen-free sample, allthe flat and deep traps are effective and the diffusivity isnot defined. The trapping effect of a steel is characte-rized by the retention of the permeation through themembrane, i.e. the ’time-lag’ till the hydrogen appears atthe exit side after the start of charging. A measure for theamount of hydrogen trapped is the charge QT

corresponding to the hatched area in Figure 3b, whichcan be obtained by the integration of the permeationcurve Ip(t) to te.

If the hydrogen concentration at the entry side is keptconstant, the trapping effect is described by:

tT = 2Q / Ip (22)

For pure trap-free iron

tL = 2Q / Ip = d2 / 6 DH (23)

where DH is the hydrogen diffusivity in iron without asignificant trapping effect.

The theory of McNabb and Foster leads to anexpression describing the retardation of hydrogendiffusion by the action of flat traps using equations (12)and (13).

Thus from the non-steady-state permeation curvesthe number of flat traps (Nft) can be determined. Thebinding energy of the hydrogen atoms in the flat trapsand their concentration can be obtained using equation

(12) for calculating Deff from measurements at differenttemperatures. (This evaluation leads to a value for ∆Hft

0,and ∆Sft

0 is assumed to be small).From time-lag measurements the effect of deep traps

can be derived using the theory of McNabb and Foster 33

which leads to the expression:

t

tT

L

−

=1

31

2 21

2

αβ β β

β αdt

dt dt dt

dt ftln+ − ⋅ +

+( )

with α dt dt dtN K / N= ⋅ βdt H dtc K / N= ⋅ (24)

Measurements of the time lag as a function of thehydrogen pressure can be fitted according to equation(24) and yield the number of deep traps and the bindingenergy, in this case ∆Gdt

0 is obtained. The value of αft forthe flat traps must be known from non-steady-statemeasurements of the permeation transients in a range ofelevated aH

40.

3 RESULTS

3.1 Effects of C, S, P, Mn, Si, Cr, Ni, Sn and Cu oncorrosion and H absorption

During the pickling of steels hydrogen is oftenabsorbed which may cause defects and failures in thematerial. The hydrogen absorption is affected bycomponents in the electrolyte, but also by thecomposition of the steel surface on which alloyingelements and impurities may be enriched after hot rollingor after some dissolution in the pickling acid. Theretarding effects of Cu on corrosion and H-absorptionhave been reported 21,22. Phosphorous is reported toenhance the corrosion and H absorption, while Habsorption 19,20,21 is also increased by S, Se, As, Sb andSn 19,23,24,25. These effects have been investigatedquantitatively for various binary Fe alloys (see table 1),using the electrochemical double cell 26,27.

Table 1: Compositions (wt%) of investigated alloys (only mainalloying addition, otherwise relatively pure, see original papers)Tabela 1: Sestava (mas.%) raziskanih zlitin (samo glavni legirnielementi, sicer ~iste zlitine, glej izvirne ~lanke)Binary alloys/Binarne zlitine 26-28

Fe-0.20 C, Fe-0.046%S, Fe-0.052%P, Fe-0.1%P, Fe-3.08%Si,Fe-1.53%Mn, Fe-5.19Cr, Fe-10.4%CrFe-5.07%Ni, Fe-10.2%Ni, Fe-0.55%Cu, Fe-3.4% Cu, Fe-0.17%SnFe-0.21%Ti, Fe-0.19%V, Fe-0.3%Zr, Fe-0.35%Nb, Fe-0.36%MoTernary alloys/Ternarne zlitine 34-40

alloysFe-Me-C

% Me % C alloysFe-Me-N

% Me % N

Fe-Ti-C 0.22 0.074 Fe-Ti-N 0.18 0.042Fe-V-C 0.19 0.081 Fe-V-N 0.20 0.048Fe-Zr-C 0.27 0.067 Fe-Zr-N 0.63 0.081Fe-Nb-C 0.35 0.066 Fe-Nb-N 0.35 0.043Fe-Mo-C 0.33 0.065 Fe-Mo-N 0.37 0.061

For the determination of the hydrogen activities andconcentrations established during corrosion the



permeation coefficient Φ° had to be measured. Figure4a shows that most of the alloying elements cause only aslight reduction in Φ°: the greatest effect was observedfor 3,1 wt%Si and 10,4%Cr. For these alloys the Hdiffusivities were also determined, see Figure 4b. Mostelements have only a small effect on the diffusivity, butCr clearly retards H diffusion, indicating trapping by Cratoms. A calculation of the hydrogen solubility ispossible from cH = Φ0 / Deff, and the data for the Fe-Sialloy render a clearly decreased solubility 26. The studiedalloying and impurity elements affect H permeation onlyweakly, Si and Cr decrease the permeation flow, Cr trapsH and thus decreases DH, and Si tends to decrease the Hsolubility.

The investigated iron alloys, see table 1, werecorroded in 1M H2SO4 and analyzed afterwards bytaking AES sputter profiles 27. Carbon and sulfur areenriched in the surface after corrosion, as is phosphorus,which according to XPS spectra is present as bothphosphide and phosphate. In contrast, silicon andchromium are depleted near the surface, these elementsare obviously dissolved preferentially. Of the nobler

elements, Ni is not clearly enriched, but minordissolution of Sn and Cu leads to the enrichment of theseelements. Current density potential curves measured onthese binary alloys in 1M H2SO4 indicate, by their shiftafter some time of corrosion, if anodic or cathodicreactions are changed by these enrichments 28, seeFigure 5. The most interesting data are the mass loss bycorrosion, and the hydrogen activity established duringcorrosion, which have been recorded after periods ofcorrosion in 1M H2SO4 at 25 °C, see Figure 5. The datashow a strong increase in corrosion by S and P, whereasNi, Sn and Cu reduce the amount of corrosion, the otherelements have minor effects. The high mass loss of theFe-Mn alloy was obviously due to its relatively high Scontent (80 wt.ppm compared to 20 wt.ppm in theunalloyed Fe). The effects of S and P are most probablycaused by the catalytic effects of sulphide and phoshideanions on the anodic dissolution of iron. Detailed studieshave been conducted on the effects of Cu 18 and of P 34.

The total hydrogen content of the samples afterdifferent periods of corrosion in 1 M H2SO4 wasdetermined by chemical analysis (hot extraction). The Hcontent of the pure Fe and the Fe-C, Fe-Ni and Fe-Cualloys are low, around the detection limit of the method(0.4·10-6 mol H/cm3). The hydrogen content is enhancedin the Fe-Cr alloys due to the trapping effect. For somealloys the H content during corrosion in 1 M H2SO4 goesthrough a maximum and then decreases, e. g. for Fe-Sncaused by increasing Sn enrichment, and for Fe-P by theformation of a phosphate film. Most striking is the highH content of the Fe-S alloy.

At high values of aH the amount of hydrogen in Feand its alloys increases more than the hydrogen activity,



Figure 4: Hydrogen permeation and diffusion in binary Fe alloys 26:a) permeation coefficient Φ° measured at pH2 = 1 bar, b) hydrogendiffusivities Deff measured after saturation of the deep trapsSlika 4: Permeacije in difuzija vodika v binarnih zlitinah `eleza 26: a)koeficient permeacije Φ°, izmerjen pri pH2 = 1 bar, b) difuzivnostvodika Deff, izmerjena po nasi~enju globokih pasti

Figure 5: Results of corrosion experiments in 1 M H2SO4 at 25 °C onbinary Fe alloys 27, mass loss (∆m) and hydrogen activity (aH) after 25hours corrosion, corrosion potential and amount of dissolved alloyingelement (related to the initial composition, not determined for Fe-0.2Cand Fe-0.046S)Slika 5: Rezultati korozijskih poskusov na binarnih Fe-zlitinah v 1 MH2SO4 pri 25 °C 27, izguba mase ∆m in aktivnost vodika po 25 urahkorozije, korozijski potencial in koli~ina raztopljenih legirnihelementov (glede na za~etno sestavo, ni dolo~ena za Fe-0,2 C inFe-0,046 S).

due to the formation of dislocations, defects and micro-cracks.

3.2 Effects of Mo, V, Nb, Ti, Zr and their carbides andnitrides on H diffusivity and solubility

Fine-grained microalloyed steels have found manyapplications because of their good mechanicalproperties, high strength and good ductility, but they aresusceptible to hydrogen embrittlement. Their micro-structure is stabilized by finely divided carbides, nitridesand carbonitrides of the elements V, Nb and Ti, whichpin grain boundaries and dislocations. These precipitatesgenerate a lot of extended sites which are suitable forhydrogen atoms, i.e. hydrogen traps. Such traps affectthe solubility and diffusivity of hydrogen. To study theeffect of the alloying elements Me used in microalloyedsteels and their carbides, binary Fe-Me and ternaryFe-Me-C and Fe-Me-N melts were prepared with about0,2 at.% Me = Mo, V, Nb, Ti, Zr, see table 1. The ingotswere forged and solution annealed at high temperature.The binary alloys were heat treated at 750 °C for 30 min,the ternary alloys for 300 h at 600 °C.

The materials were characterized by optical andelectron microscopy 35-40. In the Fe-Mo-C alloy, besidesthe coarse primary Fe3C-Mo7C3 carbides, there were fineMo2C needles in the bulk which cause strains anddislocations in their neighbourhood due to theirhexagonal structure and incoherence with the iron lattice34-36. In contrast, the fine VCx particles in Fe-V-C alloyare finely dispersed, they are cubic and semicoherentwith the Fe lattice and cause less strains. The cubic TiC,NbC and ZrC particles were less numerous and relativelycoarse, due to their low solubility. The nitrides are evenless soluble and during the solution annealing much lessdissolution of the primary nitrides takes place, so thatthere are less finely dispersed nitride precipitates thancarbide precipitates formed after the heat treatment 40.

For these alloys it was most interesting to see, howthe alloying elements and their carbides and nitridesaffect the diffusivity and solubility of hydrogen by the

various different traps generated by the presence of theelements and their precipitates.

For the binary Fe-Me alloys the steady-statepermeation measurements produced data for Φ° whichare a little lower than for pure iron, the temperature



Figure 6: Hydrogen permeation and diffusion in ternary Fe alloyswith precipitates of Ti-, V- and Nb carbides and nitrides 40: a)permeation coefficient Φ° from steady-state experiments, b) hydrogendiffusivities from non-steady-state experimentsSlika 6: Permeacija in difuzija vodika v ternarnih Fe-zlitinah z izlo~kiTi-, V- in Nb- karbidov in nitridov 40: a) koeficient permeacije Φ° zaposkuse v stacionarnem stanju, b) difuzivnost vodika iz poskusov vnestacionarnem stanju

Table 2: Data on permeation, diffusivity, solubility and trapping of H in ternary alloysTabela 2: Podatki o permeaciji, difuzivnosti, topnosti in ujetju H v ternarnih zlitinah

Alloy Φ° ⋅ 1013

mol/cm ⋅ sDeff ⋅ 106

cm2/scHeff ⋅ 108

mol/cm3cHtotal ⋅ 106

mol/cm3Nft ⋅ 103

mol/cm3-∆HftkJ/mol

Ndt ⋅ 106

mol/cm3-∆GdtkJ/mol

Fe-Ti-C 1.28 0.72 3.5 10 28 20.6 10 58.5Fe-Ti-N 2.2 6.1 4.1 1.9 3.0 20.7 1.9 60.5Fe-V-C 2.12 3.5 7.2 1.8 22 17.2 1.8 57.0Fe-V-N 2.15 2.7 9.2 3.0 14 18.9 3.1 56.0Fe-Zr-C 2.05 24 1.1 0.4 0.83 19.9 0.4 58.5Fe-Zr-N 2.09 3.0 8.5 16 7.3 20.4 17 56.1Fe-Nb-C 2.16 0.82 30 13 60 18.3 13 56.0Fe-Nb-N 2.23 5.3 4.7 3.5 9.2 18.2 3.7 54.9Fe-Mo-C 2.0 42 0.6 0.05 4.0 13.9 0.045 56.5Fe-Mo-N 1.68 0.27 0.92 0.04 0.9 19.3 0.035 56.0

Fe(60% def,) 2.51 96 26 0.72 27.9 0.62 55.5

dependency corresponded to the value ∆H = 34,3kJ/molH for pure iron. Also for the ternary Fe-Me-C andFe-Me-N alloys with carbides and nitrides, Φ° wasdecreased only by a tortuosity factor see table 2 andFigure 6a. This decrease is most probably caused by thereduction in the cross-sectional area free for diffusion,due to the precipitates, and possibly, regions withcompressive stresses. The effect on Φ° is low for most ofthe binary and ternary alloys, the largest effect is for theFe-Ti-C alloy where a high concentration of particleswith sizes 1-5 nm was observed by TEM.

In contrast, the non-steady-state measurements of Deff

showed strong effects of the alloying elements as well asof the precipitates. H diffusivity in the Fe-Me alloysdecreases strongly with increasing atomic size of theelements in the sequence V, Mo, Ti, Nb and Zr 37. Whilethe effective diffusivity decreases, the hydrogensolubility increases, both caused by flat traps andextended interstitial sites near the Me substituting atoms.In addition, attractive chemical interactions may occurbetween these hydride-forming elements and thedissolved H-atoms. The binding energies are between -5kJ/mol H for V and -21 kJ/mol H for Zr 37.

The effective diffusivity is also decreased in thepresence of the carbides and nitrides by about one order

of magnitude, see Figure 6b, and this effect is explainedby flat traps with binding energies in the range -17kJ/mol H (Fe-V-C) and -21 kJ/mol H (Fe-Ti-C,Fe-Ti-N). The amount of dissolved hydrogen is mainlydetermined by the presence of deep traps with bindingenergies between -55 kJ/mol H (Fe-Nb-N) and -60.5kJ/mol H (Fe-Ti-N). The number of flat traps is muchhigher 10-2 - 10-3 mol/cm3 compared with 10-6 - 10-5

mol/cm3 of deep traps. However, analysis of the datashows that most of the hydrogen is in the deep traps (>96%), the concentration in the flat traps is much smaller(1-4 %), and the ideally dissolved hydrogen is much less(< 0,1 %). This is valid at aH = 1, however, the amount ofmobile H in ideal solution and in flat traps increases inproportion to aH and becomes greater than the amount ofH in deep traps at the critical high values of aH > 100.

Upon corrosion in 1 M H2SO4 at 25 °C the alloyingelements from the binary Fe-Me alloys are dissolvedwithout enrichment at the surface, as observed withAES. Mo and Zr cause small reductions of cathodichydrogen formation and corrosion current. The carbidesshow some corrosion resistance and appear to beenriched in a surface oxide film and also have minoreffects on the corrosion behaviour. Mass loss andhydrogen activity established at the corrosion potential



Table 3: Compositions (wt%) of investigated steels 46

Tabela 3: Sestava (mas.%) raziskanih jekel 46

Steel C Si Mn P S Mo Ni Cr Cu V NbSt 0 0.05 0.45 1.45 0.009 0.001 0.007 0.23 0.03 0.3 0.035 0.047St 1 - 4 similar, different thermomechanical treatments (see table 4)St 5 0.6 0.28 0.83 0.008 >.001 - - 0.033St 6 0.04 0.29 1.36 0.009 0.14 0.2 - 0.034St 7 0.03 0.3 1.28 0.006 - 0.04 0.038A516 0.188 0.267 1.05 0.008 0.001 0.055 0.2 0.189 0.113 - -E355 0.116 0.275 1.34 0.008 0.0005 0.059 0.208 0.1 0.12 0.026E500 0.069 0.248 1.36 0.0009 0.0001 0.208 0.505 0.147 0.046 -

Table 4: Data on permeation. diffusion. solubility and trapping of H in low-alloy steelsTabela 4: Podatki o permeaciji, difuziji, topnosti in ujetju H v malo legiranih jeklih

Steel Φ° ⋅ 1013

mol/cm2⋅sDeff ⋅ 106

cm2/sceff ⋅ 108

mol/cm3ctotal ⋅ 108

mol/cm3Nft ⋅ 103

mol/cm3-∆HftkJ/mol

Ndt ⋅ 108mol/cm3

-∆GdtkJ/mol

St 0 TM(γ) surface 1.63 10.4 2.34 9.58 0.63 23.0 7.4 58.0centre 1.63 16 1.53 8.03 0.86 21.0 7.5 53.2

St 1 TM(γ) 1.65 21.5 1.17 4.22 1.4 19.0 3.2 56.0St 2 TM(α+γ) 1.58 18.1 1.4 5.9 2.6 18.0 5.0 54.0St 3 TM(γ) + ACC 1.58 17.6 1.42 6.22 3.1 17.6 6.0 52.0St 4 Q + T 1.79 10.2 2.44 27.3 0.46 23.9 29.0 53.0St 5 TM(γ) 1.74 39.4 0.63 6.54 8.1 12.5 6.5 54.3St 6 TM(γ) + ACC 1.56 26.4 0.94 4.76 4.1 15.6 5.0 51.5St 7 TM(α + γ) + ACC 1.81 33.2 0.75 6.1 7.5 13.3 6.0 53.8

A 516 normalized 1.51 25.1 0.99 8.2 3.6 16.1 7.5 56.5E 355 " 1.5 23.5 1.01 7.32 1.0 19.6 6.5 57.3E 500 Q + T 1.6 14.5 1.71 48.6 5.0 17.0 5.5 53.2Fe (60 % def.) 2.51 1.4 17.4 75.8 0.72 27.9 62 55.5

do not differ significantly from the values for pure iron,however, the amount of hydrogen dissolved isconsiderably enhanced by the alloying elements: in thesequence Mo < V < Nb < Ti < Zr and also for theFe-V-C alloy with the finely distributed vanadiumcarbides and the large number of traps.

3.3 Effects of microstructure in pipeline steels

Various low-alloy steels, such as those used forpipelines, were investigated in terms of hydrogenpermeation, solubility and diffusion 46, they had differentcontents of microalloying additions, V and Nb, and wereprocessed differently see tables 3 and 4.

In the first series of experiments all the investigatedsteels were charged with hydrogen from the gas phase atpH2 = 1 bar and 25 °C, the permeation current wasfollowed until it reached the steady-state permeationcurrent Ip, this measurement yields the amount Q of Habsorbed in the steel and the permeation coefficient Φ°,which was determined for the temperature range 10-80°C.

The data for the commercial pipeline steel St0 showno clear differences for samples taken from the centreand the more deformed surface region, also heattreatment at 240 °C for 80 h had no notable effect on Φ°.Furthermore, the effect of different thermomechanicaltreatments on Φ° is small, see table 4. In all cases, thedecrease of Φ° compared to pure iron can be largelyexplained by the effect of the Si content. Silicondecreases the H solubility and thus causes a decrease inΦ° (see chapter 3.1).

The effects of microalloying elements on Φ°decrease in the sequence V > Ti > Nb, they are causedby the precipitation of carbides. The small VCx particlesare very finely dispersed, VCx and TiC are precipitatescoherent with the iron matrix, both cause internalstresses and these effect a decrease of Φ°. Internalstresses also result from fast cooling after welding, thuswelding simulation also causes a decrease in Φ°. Sincethe temperature dependence of Φ° in all cases can bedescribed by the activation energy valid for pure iron34,3 kJ/molH, the effects described are just decreasingthe temperature-independent factor, and can beexplained by the reduced cross-section of permeation, asaffected by the carbides and by compressive stresses.

The diffusivity Deff, as affected by the flat traps,shows much stronger deviations from the ideal value forpure iron than the permeation coefficient see table 4. Deff

is decreased by increasing trap density, especially in therange of higher temperatures, and a higher bindingenergy causes a steeper decrease mainly in the range oflower temperatures. The effect of high binding energiesis observed to correlate with high dislocation densitiesand internal stresses. This is true for the steelsinvestigated in this study, in the case of the stronglydeformed near-surface region of the pipeline steel St 0

and after different thermomechanical treatments, and forthe heat-affected zones (HAZ), especially after veryrapid cooling from high temperatures. By tempering, thequenching effect can be partially removed, but only ifthe microstructure is not stabilized by fine carbidedispersions. Especially the finely dispersed VCx

precipitates, which are able to stabilize a dislocationnetwork and cause a relatively low value of Deff,more-or-less independent of heat treatments and weldingsimulations.

The concentration of mobile hydrogen, which iseffective in diffusion and steady-state permeation is highfor the steels with high dislocation densities and internalstresses. This is demonstrated for the samples from thesurface region of the pipeline steel, for the quenched andtempered steels (St4, E500) and for the various HAZstructures after welding simulation. The disorderedmicrostructures are stabilized again, more or less, by thecarbide formers Ti, Nb and especially V (HAZstructures). Tempering decreases the internal stressesand simultaneously the trapping effect and hydrogenconcentration cH

eff are reduced. Only the presence of Vstabilizes these data.

Information on the deep traps was obtained fromtime-lag measurements according to equation (24), fromfitting plots of (tT - tL)/tL vs. (pH2)-1/2. The trapping bydeep traps is enhanced for steels which were temperedafter a quench, i.e. quenched and tempered steels andHAZ structures. For the base steel without Ti, Nb and Vthis effect is not observed, and we may conclude that thenumber of deep traps is connected to the presence ofcarbide precipitates. Obviously, the number of deep trapsincreases with the increasing number of carbide/ferriteinterfaces and for incoherent precipitates. Titaniumcauses large effects with deep traps, due to theincoherence of TiC. Besides the effects of carbides,deformation also leads to an increased number of deeptraps.

The effect of deep traps depends on their number andon the binding energy. The latter influence is minor,since even at the hydrogen activity aH = 1, the saturationof traps with ∆G = -51 kJ/mol is already 73%, at aH = 10it is 96 %. The values found for the investigated steelsare between -51 and -58 kJ/mol H, and thus very near tothe value determined for pure iron with a highdeformation and dislocation density: -55.5 kJ/molH. Inthe literature, values of -59 to -60 kJ/molH are given fordeep traps in dislocations and grain boundaries 8.Therefore, we can conclude that deep traps at, or in theneighbourhood of, carbides are just extended interstitialsites, as in dislocation cores and grain boundaries.Special chemical interactions of H with carbides and alsowith nitrides are improbable, and high binding energiesof -90 kJ/mol H to TiC 43,44 are not confirmed by ourwork.

The total hydrogen content is the sum of thehydrogen in deep traps and the mobile hydrogen. The



latter is in ideal solution and in flat traps, and itsconcentration is approximately proportional to aH. Inpure recrystallized iron and steels with low strength, athigh aH sources of dislocations become active and newdislocations are generated 45. For the ternary alloysFe-V-C and Fe-Ti-C with maximum strength ∼ 350 MPathe trap density was increased above aH = 30, and thetotal H content determined by permeation study and byhot extraction were about equal. The formation ofdislocations and traps will need higher hydrogenactivities, aH > 70 for fine-grain steels with higherstrength. For lower values of aH the total H content canbe satisfactorily calculated from the permeation data.Such a calculation shows that at aH = 1 the hydrogen indeep traps prevails, whereas at the higher values of aH >70 the amount of diffusible H becomes higher for all theinvestigated steels.

3.4 Fracture behaviour in the presence of hydrogen

The effect of hydrogen on the fracture of theinvestigated steels (in table 3) was observed in constant-

extension-rate tests (CERT), conducted in: 1 M H2SO4 +10-3M As2O3, at charging currents -iC = 0.1 to 8 mA/cm2.The hydrogen activities could be established between aH

= 50 and 1500, as was derived from the dependence ofaH on iC which was determined before by permeationstudies. As a measure of the effect of H on themechanical properties of steels, the decrease of the workneeded for fracturing the steels was measured, i.e. therelative work of fracture

Wf = (W/Rm)(W0/Rm

0) = Wf / Wf0 (25)

where Rm is the maximum tensile strength, Rm0 without

hydrogen.At first a study was conducted on the effect of

extension rate µ in tests on specimens which were notprecharged with H. The processes leading to thehydrogen effects, H diffusion and local accumulation,H-induced crack formation and crack growth need sometime, therefore the extension rate µ must be low enoughto realize the hydrogen effects. The results indicate that arate lower than 10-5/sec should be used to obtain reliabledata for Wf, see Figure 7a. Another important result wasobtained in tests with different steels at various values ofaH, as established by varying the charging current iC. Thework of fracture Wf decreases with increasing aH, but atvalues aH > 500, Wf remains largely constant. Smalldifferences in the hydrogen susceptibility of the steelscan be seen from the data in Figure 7b.

In addition, a series of tests were performed underthree different conditions:a) CERT at aH ∼ 500, without precharging;b) CERT at aH ∼ 1000-1500, without precharging;c) CERT in air (aH = 0) after 96 hours precharging at aH

∼ 1200, and 25 h in air at ambient temperature.



Figure 7: Results of constant-extension-rate tests (CERT) withseveral pipeline steels 46 in electrolyte 1 M H2SO4 + 10-3 M As2O3under cathodic charging with hydrogen; relative fracture energydependence on a) the extension rate µ at aH = 300, b) the hydrogenactivity aHSlika 7: Rezultati poskusov s konstantno hitrostjo iztezanja (CERT)na ve~ jeklih za cevovode 46 v elektrolitu 1 M H2SO4 + 10-3 As2O3 prikatodnem nasi~enju z vodikom; relativna energija preloma vodvisnosti od: a) hitrosti iztezanja µ pri aH = 300, b) aktivnosti vodikaaH

Figure 8: Relative fracture energies of some pipeline steels fromconstant-extension-rate tests (CERT) 46 at µ = 10-5/s: a) CERT inelectrolyte at aH ≈ 500, b) CERT in electrolyte at aH ≈ 1000 - 1500, c)CERT in air after 96 h precharging and 24 h hydrogen desorption atroom temperatureSlika 8: Relativna energija preloma nekaterih jekel za cevovode,dolo~ena s poskusi s konstantno hitrostjo iztezanja (CERT) 46 pri µ =10-5/s: a) CERT v elektrolitu pri aH ≈ 500, b) CERT v elektrolitu priaH ≈ 1000-1500, c) CERT v elektrolitu po 96 urah katodnega nasi~e-nja in 24 urne desorpcije vodika pri sobni temperaturi

Series (a) and (b) gave no clear differences in Wf, seeFigure 8a, which confirms again that the fracture workbecomes independent of aH above aH ∼ 500. Certaindifferences in H susceptibility can be seen for thedifferent steels, depending on the microstructure.

Most important is the result of the last series (c),which gave values for Wf only a little smaller than theoriginal value W°f in absence of hydrogen, or even equalto W°f and higher. While the steels are held in air atroom temperature all mobile H will desorb, but thehydrogen in deep traps will stay in the steels. So theseexperiments clearly prove that it is mainly the mobile Hand not the H in deep traps which is responsible for theH effects on fracture.

A welding simulation enhances the susceptibility toH, as shown by the tests on one of the steels treated intwo different ways: 1. original state, 2. 1 sec 1250 °C +∆T(800-500 °C) = 15 sec, 3. 1 sec 1250 °C + ∆T(800-500 sec) = 15 sec + 2 hours 600 °C. Both weldingsimulations cause a significant decrease in Wf, eventempering at 600 °C does not improve the behaviour.

4 CONCLUSIONS

Knowledge on hydrogen absorption and diffusion insteels is important for understanding defects andcracking caused by hydrogen in steel. Extensive studieshave been conducted on iron, binary and ternary ironalloys and steels, using the electrochemical double-cellmethod, for determing the H permeation, diffusivity,solubility and activity upon charging the specimens fromthe atmosphere or from an electrolyte.

The effects of the alloying elements C, S, P, Mn, Si,Cr, Ni, Sn and Cu on hydrogen permeation and diffusionhave been studied on binary Fe alloys. The permeationcoefficient Φ° = cH⋅DH is considerably decreased by Siand Cr, Si reduces the solubility cH and Cr reduces thediffusivity DH by trapping. The surface layercomposition after corrosion in 1 M H2SO4 wasinvestigated by AES and XPS, Cu, Sn, P and C areenriched after corrosion, Cr is preferentially dissolved.Cu, Sn and Ni inhibit the dissolution and decrease thehydrogen activity during corrosion. S, P and Mn enhancethe corrosion attack, and increase the hydrogen activity.

The H diffusivities in Fe alloys and steels arestrongly decreased at low temperatures < 200 °C bytrapping, i.e. the binding of H to traps at an energy levellower than in the normal interstitial sites. Data ontrapping were derived according to the theory ofMcNabb and Foster 33, from non-steady state permeationstudies on binary Fe-Me, ternary Fe-Me-C and Fe-Me-Nalloys (Me = Ti, V, Zr, Nb, Mo). For all alloys, flat trapscan be found with binding energies around -19 kJ/mol,which are filled to a level more-or-less dependent on aH

and affect the diffusion of the mobile H. In contrast, thedeep traps with binding energies around -57 kJ/mol arealready filled at low aH and the hydrogen in deep traps is

virtually immobile at ambient temperature. There is noindication of a special chemical bonding of H to carbidesor nitrides, it can be assumed that the deep traps are sitesin dislocation cores, grain boundaries and at interfaces,whereas the flat traps are extended lattices sites nearsuch defects.

These considerations are also true for low-alloysteels, as shown for some pipeline steels. These steelswere tested for their hydrogen susceptibility in constant-extension-rate tests, conducted after charging withhydrogen from an electrolyte. No great differences wereobserved in the dependence on their microstructure. Thework in fracturing at a low extension rate (10-5/sec)decreased with increasing aH but was about constant at aH

> 300. After precharging and desorption of the mobilehydrogen at room temperature there was no effect of thehydrogen in deep traps, on the work of fracture. Only themobile H in an ideal solution or in flat traps is involvedin fracture, its concentration increases approximately inproportion to aH. Transgranular cracks occur in highlystressed regions after mobile H has concentratedsufficiently. The immobile H in deep traps has no effecton the fracture behaviour of the steels, thus there is nodirect correlation between total hydrogen content and thesusceptibility to hydrogen-induced cracking andstress-corrosion cracking.

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absorption and diffusion of hydrogen in steels

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